JP2005060840A - Steel pipe with low yield ratio, high strength, high toughness and superior strain age-hardening resistance, and manufacturing method therefor - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a steel pipe with the low yield ratio, high strength, high toughness and superior strain age-hardening resistance, which is manufactured at high manufacturing efficiency and at a low cost. <P>SOLUTION: The steel pipe with the low yield ratio, high strength, high toughness and superior strain age-hardening resistance comprises, by mass%, 0.03-0.1% C, 0.01-0.5% Si, 1.2-2.5% Mn, 0.08% or less Al, one or more elements selected among 0.005-0.04% Ti, 0.005-0.07% Nb and 0.005-0.1% V, while controlling a ratio, by atom%, of a C content to the total content of Ti, Nb and V, namely C/(Ti+Nb+V) to 1.2-3, and the balance substantially Fe; and has a metallographic structure of a three-phase structure consisting substantially of ferrite, bainite and island martensite, in which an area fraction of the island martensite is 3 to 20%. <P>COPYRIGHT: (C)2005,JPO&NCIPI

Description

  The present invention relates to a large-diameter welded steel pipe (UOE steel pipe, spiral steel pipe) that is suitable for a line pipe for mainly transporting crude oil and natural gas and has a small material deterioration after coating treatment, and a method for producing the same.

  In line pipes that mainly transport crude oil and natural gas, in addition to high strength and toughness, a low yield ratio is also required from the viewpoint of earthquake resistance. In general, it is known that the yield ratio of steel can be reduced by making the microstructure of steel a structure in which a hard phase such as bainite or martensite is appropriately dispersed in a soft phase such as ferrite. It has been. As a production method for obtaining a structure in which a hard phase is appropriately dispersed in such a soft phase, quenching from a two-phase region of ferrite and austenite (Q ′) between quenching (Q) and tempering (T). There is known a heat treatment method (see, for example, Patent Document 1).

  Further, as a technique for achieving a low yield ratio without performing a complicated heat treatment as disclosed in Patent Document 1, the rolling of the steel material is completed at the Ar3 transformation point or higher, and the subsequent accelerated cooling rate and cooling stop temperature. Is known to achieve a low yield ratio with a two-phase structure of acicular ferrite and martensite (see, for example, Patent Document 2).

  However, welded steel pipes such as UOE steel pipes and ERW steel pipes used for line pipes are formed by cold forming the steel sheet into a tubular shape, welding the butt, and then usually coating the outer surface of the steel pipe from the standpoint of corrosion protection. Therefore, strain aging occurs due to processing strain during pipe making and heating during coating treatment, and yield stress increases. Therefore, even if the low yield ratio of the raw steel plate is achieved by the method described above, it is difficult to achieve a low yield ratio in the steel pipe.

Strain aging is suppressed by limiting the C and N contents that cause strain aging, and adding Nb and Ti to bond with C and N as a steel material with excellent strain aging characteristics and its manufacturing method. There is a known method (see, for example, Patent Document 3).
JP-A-55-97425 Japanese Patent Laid-Open No. 1-176027 JP 2002-220634 A

  However, in the technique described in Patent Document 3, as shown in the examples, since the hot rolling finish temperature is low, the productivity is extremely lowered and the manufacturing cost is increased.

  As described above, in the conventional technique, it is difficult to manufacture a steel pipe having a low yield ratio even after the coating process without reducing productivity and without increasing manufacturing cost.

  Therefore, an object of the present invention is to solve such problems of the prior art, and to provide a low yield ratio high strength high toughness steel pipe excellent in strain aging characteristics that can be manufactured at high manufacturing efficiency and low cost, and a method for manufacturing the same. There is.

The features of the present invention for solving such problems are as follows.
(1) By mass%, C: 0.03 to 0.1%, Si: 0.01 to 0.5%, Mn: 1.2 to 2.5%, Al: 0.08% or less And at least two selected from Ti: 0.005-0.04%, Nb: 0.005-0.07%, V: 0.005-0.1%, with the balance being substantially And C / (Ti + Nb + V), which is a ratio of the amount of C in atomic% and the total amount of Ti, Nb, and V, is 1.2 to 3, and the metal structure is substantially composed of ferrite and bainite. A low yield ratio, high strength, high toughness steel pipe excellent in strain aging resistance, characterized by having a three-phase structure with island martensite and having an area fraction of island martensite of 3-20%.
(2) Further, in terms of mass%, Cu: 0.5% or less, Ni: 0.5% or less, Cr: 0.5% or less, Ca: 0.0005 to 0.003%, B: 0.005 % Low yield ratio high strength high toughness steel pipe excellent in strain aging resistance as described in (1), containing one or two or more selected from below.
(3) After the steel having the component composition described in (1) or (2) is heated to a temperature of 1000 to 1300 ° C. and hot-rolled at a rolling end temperature of Ar 3 temperature or higher, 5 ° C./s or higher Accelerated cooling to 450 to 650 ° C. at a cooling rate of 5 ° C., followed by immediate reheating to 550 to 750 ° C. at a heating rate of 0.5 ° C./s or more, and the metal structure is substantially ferrite, bainite and island As a steel plate having a three-phase structure with martensite and an island-like martensite area fraction of 3 to 20%, the steel plate is cold formed into a tubular shape, and the butt portion is welded to form a steel pipe A method for producing a low yield ratio, high strength, high toughness steel pipe excellent in strain aging characteristics.

  ADVANTAGE OF THE INVENTION According to this invention, the low yield ratio high strength high toughness steel pipe excellent in the strain aging characteristic can be manufactured with high manufacturing efficiency and low cost. For this reason, the steel pipe used for a line pipe can be manufactured stably in large quantities at low cost, and productivity and economy can be remarkably improved.

  In order to solve the above-mentioned problems, the present inventors have intensively studied a manufacturing method of a steel pipe original sheet, particularly a manufacturing process of accelerated cooling after controlled rolling and subsequent reheating, and as a result, the following (a) to (d) Obtained knowledge.

  (A) In the accelerated cooling process, during the bainite transformation, that is, in the temperature region where untransformed austenite exists, the cooling is stopped, and then the bainite transformation finish temperature (hereinafter referred to as the Bf point) is reheated to perform heating. Is a three-phase structure in which island-like martensite (hereinafter referred to as MA), which is a hard phase, is uniformly formed in the mixed phase of ferrite and bainite, and a low yield ratio is possible. Further, this MA is stable even when heated to 300 ° C. or lower during coating. MA can be easily identified by, for example, etching with a 3% nital solution (nitral: nitric alcohol solution) and then observing it by electrolytic etching. FIG. 2 is a photograph when the microstructure of a steel pipe is observed with a scanning electron microscope (SEM). MA is observed as a white floating portion, and MA is uniformly generated in a mixed structure of ferrite and bainite. The state can be confirmed. This MA is also stable after heating during coating.

  (B) At the time of ferrite transformation from untransformed austenite at the time of reheating, a composite precipitate containing at least two or more selected from Ti, Nb, and V is deposited. Moreover, the bainite transformed at the time of accelerated cooling dissolves C in a supersaturated state, and the above-described fine precipitates are precipitated from the bainite during air cooling after reheating. The precipitation of fine precipitates reduces the solid solution C and N that cause strain aging, so that it is possible to suppress an increase in yield stress due to strain aging after steel pipe forming and coating treatment.

  (C) Since the above composite precipitates are very fine, in addition to strengthening by bainite transformation during accelerated cooling, precipitation strengthening by fine precipitates can be obtained, so even in low-component steels with few alloying elements High strength can be achieved.

  (D) The effects of the above (a), (b) and (c) are achieved by adding a hardenability improving element such as Mn to promote the formation of MA, and at the same time, forming carbide forming elements such as Ti, Nb and V. It can be obtained by using added steel.

  The present invention was obtained from the above findings, and a bainite phase produced by accelerated cooling after rolling, and a ferrite phase and a bainite phase in which precipitates containing Ti, Nb, and V produced by subsequent reheating were dispersed and precipitated. And a low yield ratio, high strength, high toughness steel pipe having a three-phase structure in which MA, which is a hard phase, is uniformly formed and having excellent strain aging resistance.

  Hereinafter, the high-strength steel pipe of the present invention will be described in detail. First, the structure of the high-strength steel pipe of the present invention will be described.

  In the present invention, a structure in which MA, which is a hard phase, is uniformly formed in ferrite and bainite, and fine precipitates are precipitated in the ferrite phase to reduce the solid solution C and N that cause strain aging. A low yield ratio is achieved in the treated steel pipe. The mechanism of MA generation in the present invention is as follows. After heating the slab, rolling is finished in the austenite region, and then accelerated cooling is started at the Ar3 transformation temperature or higher. This is a manufacturing process in which accelerated cooling is terminated during bainite transformation, that is, in a temperature range where untransformed austenite exists, and then reheated at a temperature equal to or higher than the bainite transformation finish temperature (Bf point) and then cooled. The changes in the organization are as follows. Microstructures at the end of accelerated cooling are bainite and untransformed austenite, and reheating at the Bf point or higher causes ferrite transformation from untransformed austenite. However, since ferrite has a small amount of C solid solution, C is untransformed. Discharged into austenite. Therefore, the amount of C in untransformed austenite increases with the progress of ferrite transformation during reheating. At this time, if the hardenability is increased and austenite stabilizing elements such as Mn, Cu, Ni, etc. are contained in a certain amount or more, untransformed austenite in which C is concentrated remains even after reheating, and after reheating It is transformed into MA by cooling, and finally becomes a three-phase structure of bainite, ferrite, and MA. In the present invention, after accelerated cooling, it is important to perform reheating from a temperature range in which untransformed austenite exists, and when the reheating start temperature falls below the Bf point, bainite transformation is completed and untransformed austenite does not exist. The reheating start needs to be higher than the Bf point. Further, the cooling after reheating is not particularly specified because it does not affect the transformation of MA and the coarsening of fine carbide described later, but basically it is preferably air cooling. In the present invention, accelerated cooling is stopped in the middle of bainite transformation, and then reheating is performed continuously, so that MA that is a hard phase can be generated without lowering the production efficiency, and a composite containing a hard phase. A low yield ratio can be achieved by using a three-phase structure. The proportion of MA in the three-phase structure is preferably an area fraction of MA (ratio of the area of MA in an arbitrary cross section of the steel sheet in the rolling direction, the sheet width direction, etc.) and 3 to 20%. If the area fraction of MA is less than 3%, it may be insufficient to achieve a low yield ratio, and if it exceeds 20%, the base material toughness may be deteriorated. Further, from the viewpoint of lowering the yield ratio and the base material toughness, the area fraction of MA is particularly preferably 5 to 15%. Note that the area fraction of MA can be obtained, for example, by determining the area ratio occupied by MA by image processing the microstructure obtained by SEM observation. Further, if the MA is coarse, it becomes a starting point of fracture and deteriorates the toughness of the base material. Therefore, the average particle size of the MA is preferably 10 μm or less. The average particle diameter of MA can be obtained as an average value of the diameters obtained by subjecting the microstructure obtained by SEM observation to image processing, obtaining the diameter of a circle having the same area as each MA, and obtaining the diameter of each MA. .

  In addition, in order to suppress the yield stress increase due to strain aging after steel pipe forming and coating treatment and achieve high strength, transformation strengthening by bainite transformation during accelerated cooling and reheating after accelerated cooling into the ferrite By combining and utilizing precipitation strengthening by precipitation of fine composite carbides that precipitate, high strength can be achieved without adding a large amount of alloying elements. Ferrite is rich in ductility and is generally soft, but in the present invention, the strength is increased by the precipitation of fine composite carbide described below. When a large amount of the alloy element is not added, the strength is insufficient only by the bainite single-phase structure obtained by accelerated cooling, but the precipitation strengthened ferrite has sufficient strength. A steel plate utilizing precipitation strengthening generally has a high yield ratio, but in the present invention, low yield is achieved by uniformly generating MA having a large hardness difference from other phases such as ferrite and bainite. Furthermore, since solid solution C and N which are the cause of strain aging are fixed as fine precipitates, it is possible to suppress strain aging after heating during steel pipe forming and coating.

  Note that the metal structure substantially consists of a three-phase structure of ferrite, bainite, and island martensite, as long as the effects of the present invention are not lost, those containing structures other than ferrite, bainite, and MA, It is meant to be included in the scope of the present invention.

  When one or more different metal structures such as pearlite are mixed in the three-phase structure of ferrite, bainite, and MA, the strength decreases, so the area fraction of the structure other than ferrite, bainite, and MA is small. Moderately good. However, since the influence can be ignored when the area fraction of the structure other than ferrite, bainite and MA is low, other metal structures of less than 3% in total area fraction, that is, one or two of pearlite, cementite, etc. You may contain more than a seed. Further, it is desirable that the area fraction of ferrite is 5% or more from the viewpoint of securing strength, and the area fraction of bainite is 10% or more from the viewpoint of securing toughness of the base material.

  Next, the fine precipitate which precipitates in the ferrite phase will be described.

  In the steel pipe of the present invention, a composite precipitate containing two or more selected from Ti, Nb, and V in a ferrite phase and a bainite phase is utilized for precipitation strengthening and strain aging resistance improvement. Ti, Nb, and V are elements that form carbides in steel, and it has been conventionally performed to strengthen steel by precipitation of individual carbides. In the present invention, two kinds selected from Ti, Nb, and V are used. Compared with the precipitation strengthening by individual carbides, by adding the above, and finely dispersing and precipitating composite carbide containing two or more selected from Ti, Nb, and V in steel. It is characterized in that a large strength improvement effect can be obtained. This unprecedented strength improvement effect is due to the fact that this composite carbide is stable and has a slow growth rate, so that extremely fine precipitates having a particle size of less than 10 nm can be obtained. The number ratio of fine precipitates of the composite carbide is preferably 95% or more of the total precipitates excluding TiN. The average particle size of the fine composite carbide precipitates is obtained by subjecting a photograph taken with a transmission electron microscope (TEM) to image processing, and obtaining the diameter of a circle having the same area as each precipitate for each composite carbide. , And can be obtained as an average value of their diameters.

  Precipitates in the ferrite phase precipitate at the ferrite transformation interface from untransformed austenite during reheating, and precipitates in the bainite phase precipitate during air cooling after reheating.

  As described above, the steel pipe of the present invention has a composite structure composed of three phases of ferrite, bainite phase, and MA in which precipitates are finely precipitated. Such a structure uses steel having the following composition. It can be obtained by the following method.

  First, chemical components of the high-strength steel pipe of the present invention will be described. In the following description, all units represented by% are mass%.

  C: Set to 0.03 to 0.1%. C contributes to precipitation strengthening as a carbide and is an important element for MA formation. However, if it is less than 0.03%, it is insufficient for formation of MA, and sufficient strength cannot be secured. Addition exceeding 0.1% not only deteriorates the HAZ toughness but also deteriorates the strain aging resistance, so the C content is specified to be 0.03 to 0.1%. More preferably, it is 0.03 to 0.08%.

  Si: 0.01 to 0.5%. Si is added for deoxidation, but if it is less than 0.01%, the deoxidation effect is not sufficient, and if it exceeds 0.5%, not only the toughness and weldability are deteriorated, but also the strain aging resistance is lowered. The Si content is specified to be 0.01 to 0.5%. More preferably, it is 0.01 to 0.3%.

  Mn: 1.2 to 2.5%. Mn is added to improve strength and toughness, further improve hardenability and promote MA formation. However, if it is less than 1.2%, its effect is not sufficient, and if it exceeds 2.5%, toughness and weldability deteriorate. Therefore, the Mn content is specified to be 1.2 to 2.5%. In the case where Cu, Ni, or the like is not added in combination, it is desirable to add 1.5% or more in order to stably produce MA regardless of changes in components and manufacturing conditions.

  Al: 0.08% or less. Al is added as a deoxidizer, but if it exceeds 0.08%, the cleanliness of the steel decreases and the toughness deteriorates, so the Al content is specified to be 0.08% or less. Preferably, the content is 0.01 to 0.08%.

  The steel plate of this invention contains 2 or more types chosen from Ti, Nb, and V.

  Ti: 0.005 to 0.04%. By adding 0.005% or more, fine composite carbide is formed together with Nb and / or V, and greatly contributes to strength increase and strain aging resistance. However, since addition exceeding 0.04% causes deterioration of the weld heat affected zone toughness, the Ti content is specified to be 0.005 to 0.04%. Furthermore, when the Ti content is less than 0.02%, more excellent toughness is exhibited, so that the Ti content is preferably less than 0.005 to 0.02%.

  Nb: 0.005 to 0.07%. Nb improves toughness by refining the structure, but forms a composite precipitate with Ti and / or V and contributes to an increase in strength. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.07%, the toughness of the weld heat-affected zone deteriorates, so the Nb content is specified to be 0.005 to 0.07%.

  V: Set to 0.005 to 0.1%. V, like Ti and Nb, forms a composite precipitate with Ti and / or Nb and contributes to an increase in strength. However, if it is less than 0.005%, there is no effect, and if it exceeds 0.1%, the toughness of the weld heat affected zone deteriorates, so the V content is specified to be 0.005 to 0.1%.

The high-strength steel pipe of the present invention can obtain fine composite precipitates containing any two or more of Ti, Nb, and V by using the steel of the above components. In order to generate the above, it is necessary to limit the content ratio of the elements forming the composite precipitate as follows. That is, C / (Ti + Nb + V), which is a ratio of the amount of C in atomic% and the total amount of Ti, Nb, and V, is set to 1.2-3. Strengthening according to the present invention is due to the precipitation of fine carbides containing any two or more of Ti, Nb, and V. At this time, when the value of C / (Ti + Nb + V) represented by the atomic% content of each element is less than 1.2, all of C is consumed in the fine composite precipitate, and MA is not generated, so the yield ratio is reduced. Cannot be achieved. When C exceeds 3, C is excessive, and a hardened structure such as island martensite is formed in the weld heat affected zone, resulting in deterioration of the weld heat affected zone toughness. Therefore, the value of C / (Ti + Nb + V) is 1. 2-3. More preferably, it is 1.4-3. In addition, when content of mass% is used, each element symbol is defined as content of each element in mass% (C / 12.01) / (Ti / 47.9 + Nb / 92.91 + V / 50.94). Is set to 1.2-3.
More preferably, it is 1.4-3.

  In the present invention, for the purpose of further improving the strength toughness of the steel pipe and improving the hardenability and promoting the production of MA, it contains one or more of Cu, Ni, Cr, B, and Ca shown below. May be.

  Cu: 0.5% or less. Cu is an element effective for improving toughness and increasing strength. In order to acquire the effect, it is preferable to add 0.1% or more. However, if it is added in a large amount, weldability deteriorates.

  Ni: 0.5% or less. Ni is an element effective for improving toughness and increasing strength. In order to obtain the effect, it is preferable to add 0.1% or more, but adding a large amount is disadvantageous in terms of cost, and the weld heat affected zone toughness deteriorates. Is the upper limit.

  Cr: 0.5% or less. Cr, like Mn, is an element effective for obtaining sufficient strength even at low C. In order to acquire the effect, it is preferable to add 0.1% or more. However, if it is added in a large amount, weldability deteriorates.

  B: Set to 0.005% or less. B is an element contributing to strength increase and HAZ toughness improvement. In order to obtain the effect, it is preferable to add 0.0005% or more, but if added over 0.005%, the weldability is deteriorated, so when added, the content is made 0.005% or less.

  Ca: 0.0005 to 0.003%. Ca improves the toughness by controlling the form of sulfide inclusions. The effect appears at 0.0005% or more, and when it exceeds 0.003%, the effect is saturated, and conversely, the cleanliness is lowered and the toughness is deteriorated. And

  N: Preferably it is 0.007% or less. N is treated as an unavoidable impurity, but if over 0.007%, the weld heat affected zone toughness deteriorates, so the content is preferably made 0.007% or less.

  Furthermore, by optimizing Ti / N, which is the ratio of Ti amount to N amount, the austenite coarsening of the weld heat affected zone can be suppressed by TiN particles, and good weld heat affected zone toughness can be obtained. Therefore, Ti / N is preferably 2 to 8, and more preferably 2 to 5.

  The remainder other than the above consists essentially of Fe, and it is possible to add a trace amount of elements that do not impair the effects of the present invention, including inevitable impurities. For example, each of Mg, REM, W, and Zr may be added by 0.02% or less.

  Next, the manufacturing method of the high intensity | strength steel pipe original plate of this invention is demonstrated.

  The high-strength steel pipe sheet of the present invention uses steel having the above composition, and is hot-rolled at a heating temperature of 1000 to 1300 ° C. and a rolling end temperature of Ar 3 temperature or higher, and then at a cooling rate of 5 ° C./s or higher. Accelerated cooling to 450 to 650 ° C., and then reheating from a temperature above the Bf point to a temperature of 550 to 750 ° C. at a temperature rising rate of 0.5 ° C./s or more, the metal structure becomes ferrite and bainite. With a three-phase structure of MA, fine composite carbide containing two or more selected from Ti, Nb, and V can be dispersed and precipitated in ferrite. Here, the heating temperature, the rolling end temperature, the cooling end temperature, the reheating temperature, and the like are the average temperature of the steel sheet. The average temperature is obtained by calculation based on the surface temperature of the slab or steel plate, taking into account parameters such as plate thickness and thermal conductivity. The cooling rate is an average cooling rate obtained by dividing the temperature difference required for cooling to the cooling end temperature (450 to 650 ° C.) by the time required for the cooling after the hot rolling is completed. The temperature increase rate is an average temperature increase rate obtained by dividing the temperature difference required for reheating up to the reheating temperature (550 to 750 ° C.) by the time required for reheating after cooling. Hereinafter, each manufacturing condition will be described in detail.

  Heating temperature: 1000-1300 ° C. If the heating temperature is less than 1000 ° C, the solid solution of the carbide is insufficient and the required strength and yield ratio cannot be obtained, and if it exceeds 1300 ° C, the base metal toughness deteriorates, so the temperature is set to 1000 to 1300 ° C.

  Rolling end temperature: Ar3 temperature or higher. If the rolling end temperature is lower than the Ar3 temperature, the subsequent ferrite transformation rate is lowered, so that sufficient precipitation of fine precipitates cannot be obtained during ferrite transformation by reheating, and the strength is lowered. Moreover, since the concentration of C into untransformed austenite at the time of reheating is insufficient and MA is not generated, the rolling end temperature is set to the Ar3 temperature or higher.

  Immediately after the end of rolling, it is cooled at a cooling rate of 5 ° C./s or more. When the cooling rate is less than 5 ° C./s, pearlite is generated at the time of cooling, so that MA is not generated and strengthening by bainite cannot be obtained, so that sufficient strength cannot be obtained. Therefore, the cooling rate after the end of rolling is specified to be 5 ° C./s or more. In addition, when the cooling start temperature is lower than the Ar3 temperature and ferrite is generated, fine precipitates are not dispersed and precipitated at the time of reheating, resulting in insufficient strength and no MA formation. Therefore, the cooling start temperature is set to be higher than the Ar3 temperature. To do. About the cooling method at this time, it is possible to use arbitrary cooling equipment by a manufacturing process. In the present invention, the ferrite transformation can be completed without maintaining the temperature during the subsequent reheating by supercooling to the bainite transformation region by accelerated cooling.

  Cooling stop temperature: 450 to 650 ° C. This process is an important manufacturing condition in the present invention. In the present invention, C-concentrated untransformed austenite present after reheating is transformed into MA upon subsequent air cooling. That is, it is necessary to stop the cooling in a temperature range where untransformed austenite during the bainite transformation exists. If the cooling stop temperature is less than 450 ° C., the bainite transformation is completed, so MA is not generated during air cooling, and a low yield ratio cannot be achieved. When the temperature exceeds 650 ° C., pearlite is precipitated during cooling, so that the precipitation of fine carbides is insufficient and sufficient strength cannot be obtained. Further, C is consumed in the pearlite and MA is not generated. Specified at 650 ° C. From a viewpoint of MA production | generation, Preferably it is 500-650 degreeC, More preferably, it is 530-650 degreeC.

  Immediately after the accelerated cooling is stopped, reheating is performed to a temperature of 550 to 750 ° C. at a temperature rising rate of 0.5 ° C./s or more. This process is also an important production condition in the present invention. Fine precipitates that contribute to the strengthening of the soft phase and the improvement of strain aging characteristics are precipitated during reheating. Furthermore, due to the C transformation from untransformed austenite during reheating to ferrite transformation and accompanying untransformed austenite, untransformed austenite enriched in C during air cooling after reheating transforms to MA. In order to obtain such fine composite carbide precipitates and MA, it is necessary to reheat from a temperature above the Bf point to a temperature range of 550 to 700 ° C. after accelerated cooling. If the rate of temperature increase is less than 0.5 ° C./s, it takes a long time to reach the target reheating temperature, so that the production efficiency is deteriorated and pearlite transformation occurs, so that dispersion precipitation of fine composite carbide and MA are obtained. Thus, sufficient strength and low yield ratio cannot be obtained. If the reheating temperature is less than 550 ° C., sufficient precipitation driving force cannot be obtained, and the amount of fine composite carbide is small. This results in a decrease in strain aging characteristics and insufficient strength, and when it exceeds 750 ° C., the precipitate becomes coarse and sufficient. Since strength cannot be obtained, the temperature range of reheating is specified to be 550 to 750 ° C. In the present invention, after accelerated cooling, it is important to perform reheating from a temperature range in which untransformed austenite exists, and when the reheating start temperature falls below the Bf point, bainite transformation is completed and untransformed austenite does not exist. The reheating start needs to be higher than the Bf point. In order to reliably transform the ferrite, it is desirable to raise the temperature by 50 ° C. or more from the reheating start temperature. There is no need to set the temperature holding time at the reheating temperature. Even if it cools immediately after reheating if the manufacturing method of this invention is used, since sufficient fine composite carbide | carbonized_material will be obtained, high intensity | strength will be obtained. However, in order to ensure a sufficient fine composite carbide, the temperature can be maintained within 30 minutes. If the temperature is maintained for more than 30 minutes, the composite carbide may become coarse and the strength may decrease. Further, in the cooling process after reheating, the fine composite carbide does not become coarse regardless of the cooling rate. Therefore, it is preferable that the cooling rate after reheating is basically air cooling.

  As equipment for performing reheating after accelerated cooling, a heating device can be installed downstream of the cooling equipment for performing accelerated cooling. As the heating device, it is preferable to use a gas combustion furnace or induction heating device capable of rapid heating of the steel sheet.

  An example of equipment for carrying out the production method of the present invention is shown in FIG. As shown in FIG. 1, a hot rolling mill 3, an acceleration cooling device 4, an induction heating device 5, and a hot leveler 6 are arranged in the rolling line 1 from the upstream side toward the downstream side. By installing the induction heating device 5 or other heat treatment device on the same line as the hot rolling mill 3 that is a rolling facility and the accelerated cooling device 4 that is a subsequent cooling facility, reheating is performed quickly after the end of rolling and cooling. Since it can process, it can heat, without reducing the steel plate temperature after rolling cooling too much.

  Furthermore, the manufacturing method of a welded steel pipe is demonstrated.

  The welded steel pipe of the present invention is a steel pipe manufactured by cold forming a steel plate manufactured under the above-described manufacturing conditions into a tubular shape and welding the butted portion. The method for forming into a tubular shape is not particularly specified. The coating treatment of the formed steel pipe is preferably performed at a temperature of the steel pipe of 300 ° C. or less. When the heating temperature of the steel pipe at the time of coating exceeds 300 ° C., the strain resistance characteristics may be lowered or the stress ratio may be increased due to MA decomposition.

  FIG. 2 shows a scanning electron of a steel pipe of the present invention (0.05 mass% C-1.8 mass% Mn-0.01 mass% Ti-0.04 mass% Nb-0.05 mass% V) manufactured using the above manufacturing method. The photograph observed with the microscope (SEM) is shown. According to FIG. 2, it can be confirmed that island martensite (MA) is uniformly formed in the mixed structure of ferrite (F) and bainite (B).

  Steel of chemical composition shown in Table 1 (steel types A to I) was made into a slab by a continuous casting method, and a steel plate having a plate thickness of 18 and 26 mm was produced using this slab, and a welded steel pipe having an outer diameter of 24 and 48 inches (No. 1 To 14).

  After the heated slab is rolled by hot rolling, it is immediately cooled using a water-cooled accelerated cooling facility, reheated using an induction heating furnace or a gas combustion furnace, and a steel plate is produced. A welded steel pipe was produced by the process, and then the outer surface of the steel pipe was coated. The induction furnace was installed on the same line as the accelerated cooling equipment. Table 2 shows the production conditions of each steel pipe (No. 1 to 14). The heating temperature, rolling end temperature, cooling stop (end) temperature, reheating temperature, and other temperatures were the average temperature of the steel sheet. The average temperature was determined from the surface temperature of the slab or steel plate by parameters and calculations such as plate thickness and thermal conductivity. The cooling rate is an average cooling rate obtained by dividing the temperature difference required for cooling to the cooling stop (end) temperature after the hot rolling is finished by the time required for the cooling. The reheating rate (temperature increase rate) is an average temperature increase rate divided by the time required to reheat the temperature difference necessary for reheating up to the reheating temperature after cooling.

  The tensile properties of the steel pipe manufactured as described above were measured. The measurement results are also shown in Table 2. Tensile properties were evaluated by averaging two samples of two full thickness tensile test pieces in the vertical direction of rolling, performing tensile tests before and after coating, measuring tensile strength and yield ratio. The tensile strength of 580 MPa or more was determined as the strength required for the present invention, and the yield ratio of 85% or less was determined as the yield ratio required for the present invention.

  For base metal toughness, three full-size Charpy V-notch test pieces in the vertical direction of rolling were sampled, Charpy test was performed, the absorbed energy at −10 ° C. was measured, and the average value was obtained. The absorption energy at −10 ° C. was determined to be 200 J or more.

  For the weld heat affected zone (HAZ) toughness, a Charpy test was conducted by collecting test pieces as shown in FIG. FIG. 3 is a schematic view of a cross section of a seam welded portion of a steel pipe. From the center of the plate thickness of the seam welded portion, the notch 9 portion has a length ratio such that weld metal: HAZ = 1: 1. Three full-size Charpy V-notch test specimens 10 were collected, the Charpy absorbed energy at −10 ° C. was measured, and the average value was obtained. Those having Charpy absorbed energy at −10 ° C. of 100 J or more were considered good.

  In Table 2, all of Nos. 1 to 7 which are examples of the present invention are within the scope of the present invention in terms of chemical composition and production method, have a high tensile strength of 580 MPa or higher, and a yield ratio before coating treatment of 80%. Below, even after the coating treatment, the yield ratio was a low yield ratio of 85% or less, the strain aging resistance was excellent, and the toughness of the base material and the weld heat affected zone was good at 100 J or more. The structure of the steel pipe was a three-phase structure of ferrite, bainite, and island martensite, and the area fraction of island martensite was in the range of 3 to 20%. In addition, the area fraction of island martensite was calculated | required by image processing from the microstructure observed with the scanning electron microscope (SEM). Further, as a result of observation by transmission electron microscope and analysis by energy dispersive X-ray spectroscopy, dispersion of fine composite carbide having a particle diameter of less than 10 nm containing two or more selected from Ti, Nb and V in the ferrite phase Precipitation was observed.

  In Nos. 8 to 10, the chemical components were within the scope of the present invention, but the manufacturing method was outside the scope of the present invention, so the strength and yield ratio were insufficient. Nos. 11 to 14 had chemical components outside the scope of the present invention, so that sufficient strength was not obtained, yield ratio was high, or HAZ toughness was inferior.

Schematic which shows an example of the manufacturing line for enforcing the manufacturing method of this invention. The photograph which observed the steel pipe of the present invention with the scanning electron microscope (SEM). The schematic of the cross section of the seam weld part of a steel pipe which shows the sampling position of a test piece.

Explanation of symbols

DESCRIPTION OF SYMBOLS 1 Rolling line 2 Steel plate 3 Hot rolling mill 4 Accelerated cooling device 5 Induction heating device 6 Hot leveler 7 Steel plate 8 Weld metal 9 Notch 10 Test piece 11 HAZ
F Ferrite B Bainite MA Island martensite

Claims (3)

  In mass%, C: 0.03 to 0.1%, Si: 0.01 to 0.5%, Mn: 1.2 to 2.5%, Al: 0.08% or less, Ti: Contains at least two selected from 0.005 to 0.04%, Nb: 0.005 to 0.07%, V: 0.005 to 0.1%, with the balance being substantially Fe C / (Ti + Nb + V), which is a ratio of the amount of C in atomic% and the total amount of Ti, Nb, and V, is 1.2 to 3, and the metal structure is substantially ferrite, bainite, and island martensite. A low-yield-ratio, high-strength, high-toughness steel pipe excellent in strain aging resistance, characterized in that the area fraction of island martensite is 3 to 20%.   Further, in mass%, Cu: 0.5% or less, Ni: 0.5% or less, Cr: 0.5% or less, Ca: 0.0005 to 0.003%, B: 0.005% or less The low yield ratio high strength high toughness steel pipe excellent in strain aging resistance according to claim 1, comprising one or more selected from the group consisting of:   The steel having the component composition according to claim 1 or 2 is heated to a temperature of 1000 to 1300 ° C, hot-rolled at a rolling end temperature of Ar3 temperature or higher, and then at a cooling rate of 5 ° C / s or higher. Accelerated cooling is performed to 450 to 650 ° C., and then immediately reheated to 550 to 750 ° C. at a heating rate of 0.5 ° C./s or more, so that the metal structure is substantially composed of ferrite, bainite, and island martensite. A steel plate having a three-phase structure and an island-like martensite area fraction of 3 to 20% is formed into a tubular shape by cold forming the steel plate, and the butt portion is welded to form a steel pipe. A method for producing a low yield ratio high strength high toughness steel pipe excellent in strain aging resistance.